1. Field of the Invention
This invention relates to a method for preparing rare earth permanent magnets.
2. Prior Art
Rare earth magnets of high performance, typically powder metallurgical Sm--Co base magnets having an energy product of 32 MGOe have been produced on a large commercial scale. However, these magnets suffer from a problem that the raw materials, Sm and Co, cost much. Of rare earth elements, some elements of low atomic weight, e.g., Ce, Pr, and Nd are available in more plenty and less expensive than Sm. Iron is less expensive than cobalt. For these reasons, R-T-B base magnets (wherein R stands for a rare earth element and T stands for Fe or Fe plus Co) such as Nd--Fe--B and Nd--Fe--Co--B magnets were recently developed. One example is a sintered magnet as set forth in Japanese Patent Application Kokai (JP-A) No. 59-46008. Sintered magnets may be produced by applying a conventional powder metallurgical process for Sm--Co systems (melting.fwdarw.master alloy ingot casting.fwdarw.ingot crushing.fwdarw.fine pulverization.fwdarw.compacting.fwdarw.sintering.fwdarw.magnet), and excellent magnetic properties are readily available.
Generally, a master alloy ingot produced by casting has a structure wherein crystal grains made up of a ferromagnetic R.sub.2 Fe.sub.14 B phase (referred to as a primary phase, hereinafter) are covered with a non-magnetic R-rich phase (referred to as a grain boundary phase, hereinafter). The master alloy ingot is then pulverized or otherwise reduced to a particle diameter smaller than the crystal grain diameter, offering a magnet powder. The grain boundary phase has a function to promote sintering by converting into a liquid phase and plays an important role for the sintered magnet to generate coercivity.
One typical method for the preparation of R-T-B sintered magnets is known as a two alloy route. The two alloy route is by mixing two alloy powders of different compositions and sintering the mixture, thereby improving magnetic properties and corrosion resistance. A variety of proposals have been made on the two alloy route. All these proposals use an alloy powder having approximately the same composition (R.sub.2 T.sub.14 B ) as the primary phase of the final magnet and add a subordinate alloy powder thereto. The known subordinate alloys used heretofore include R rich alloys having a higher R content and a lower melting point than the primary phase (JP-A 4-338607 and U.S. Pat. No. 5,281,250 or JP-A 5-105915), R.sub.2 T.sub.14 B alloys containing a different type of R from the primary phase (JP-A 61-81603), and alloys containing an intermetallic compound of R (JP-A 521219).
One of the alloys used in these two alloy methods is a primary alloy of the composition R.sub.2 T.sub.14 B. If the primary alloy is produced by a melt casting process, a soft magnetic .alpha.-Fe phase precipitates to adversely affect high magnetic properties. It is then necessary to carry out solution treatment, typically at about 900.degree. C. or higher for one hour or longer. In JP-A 5-21219, for example, an R.sub.2 T.sub.14 B alloy prepared by a high-frequency melting process is subject to solution treatment at 1070.degree. C. for 20 hours. Because of such a need for high temperature, long time solution treatment, the melt casting method is against low cost manufacture. U.S. Pat. No. 5,281,250 produces an R.sub.2 T.sub.14 B alloy by a direct reduction and diffusion process, which alloy has an isometric crystal system and poor magnetic properties. A higher calcium content also precludes manufacture of high performance magnets. JP-A 4-338607 uses a crystalline or amorphous R.sub.2 T.sub.14 B alloy powder which is produced by a single roll process so as to have microcrystalline grains of up to 10 .mu.m. It is not described that the grains are columnar. It is rather presumed that the grains are isometric because magnetic properties are low. JP-A 4-338607 describes that the grain size is limited to 10 .mu.m or less in order to prevent precipitation of soft magnetic phases such as .alpha.-Fe.
With respect to thermal stability, R-T-B magnets are less stable than the Sm-Co magnets. For example, the R-T-B magnets have a differential coercivity .DELTA.iHc/.DELTA.T as great as -0.60.degree. to -0.55%/.degree. C. in the range between room temperature and 180.degree. C. and undergo a significant, irreversible demagnetization upon exposure to elevated temperatures. Therefore, the R-T-B magnets are rather impractical when it is desired to apply them to equipment intended for high temperature environment service, for example, electric and electronic devices in automobiles.
For reducing the irreversible demagnetization upon heating of R-T-B magnets, JP-A 62-165305 proposes to substitute Dy for part of Nd and Co for part of Fe. However, it is impossible to achieve a substantial reduction of .DELTA.iHc/.DELTA.T by merely adding Dy and Co. Larger amounts of Dy substituted sacrifice maximum energy product (BH)max.
JP-A 64-7503 proposes to improve thermal stability by adding gallium (Ga) while IEEE Trans. Magn. MAG-26 (1990), 1960 proposes to improve thermal stability by adding molybdenum (Mo) and vanadium (V). The addition of Ga, Mo and V is effective for improving thermal stability, but sacrifices maximum energy product.
We proposed to add tin (Sn) and aluminum (Al) for improving thermal stability with a minimal loss of maximum energy product (JP-A 3-236202). Since the addition of Sn, however, still has a tendency of lowering maximum energy product, the amount of Sn added should desirably be limited to a minimal effective level.
It was also reported to add tin (Sn) to magnets using a so-called two alloy route. The two alloy route is by mixing two alloy powders of different compositions, typically an alloy powder having a composition approximate to the primary phase composition and a subordinate alloy powder having a composition approximate to the grain boundary phase composition and sintering the mixture. For instance, Proc. 11th Inter. Workshop on Rare-Earth Magnets and their Applications, Pittsburgh, 1990, p. 313 discloses that a sintered magnet is prepared by mixing Nd.sub.14.5 Dy.sub.1.5 Fe.sub.75 AlB.sub.8 alloy powder with up to 2.5% by weight of Fe.sub.2 Sn or CoSn powder, followed by sintering. It is reported that this sintered magnet has a Nd.sub.6 Fe.sub.13 Sn phase precipitated in the grain boundary phase and is improved in thermal dependency of coercivity.
Making a follow-up experiment, we found that the Fe.sub.2 Sn or CoSn material is unlikely to fracture and thus difficult to comminute into a microparticulate powder having a consistent particle size. Then sintered magnets resulting from a mixture of an R-T-B alloy powder and a Fe.sub.2 Sn or CoSn powder contain unevenly distributed Nd.sub.6 Fe.sub.13 Sn phase of varying size. This is also evident from FIG. 5 of the above-referred article. It is thus difficult to provide thermal stability in a consistent manner. Where tin is added in the form of Fe.sub.2 Sn or CoSn powder, R and Fe in the primary phase are consumed to form Nd.sub.6 Fe.sub.13 Sn , which can alter the composition of the primary phase, deteriorating magnetic properties.